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Niobium—Base alloys
hit. J. of Refracto~ Metals & Hard Materials 12 (1993-1994) 3(13-314 © 1994 Elsevier Science Limited Printed in Great Britain. All rights reserved 0263-4368/94/$7.00 ELSEVIER
Niobium--Base Alloys E. N. Sheftel & O. A. Bannykh Baikov Institute of Metallurgy, Russian Academy of Sciences, Moscow, Russia (Received 25 March 1994; accepted 16 May 1994) Abstract: On the basis of the research carried out by the authors and the analysis of the published data, the scientific principles of the alloying and designing of structure of niobium alloys, providing high strength and workability of the alloys intended for two operating temperatures:from 1100 up to 1400°C and from 600 to 800°C, have been considered. The developing principles are based on the revealing of interdependence between phase diagrams, thermodynamiccharacteristicsof elementsand phases formingin the proper alloys as well as their structures and mechanical properties. Main attention has been paid to the alloysof Nb-Mew-X (X = O, N, C) systemswith oxide, nitride and carbide dispersoids.
INTRODUCTION
and, connected with it, the least sensibility of ductile properties to their content, it was considered as the most perspective basis for structural alloys to be used under high-temperature conditions. At that period the principal problems for Nballoys to be solved have been in searching for ways to increase high-temperature strength at satisfactory workability, as well as methods of protection against oxidisation at operating temperatures exceeding 1000°C. As a result, at the end of the 1960s and beginning of the 1970s, a great number of alloy compositions were developed which may be conventionally divided into low-, medium- and high-strength alloys that have been used for a long time operating at 1000-1200°C and a short time working at 1200-1700°C (Table 1). t,2 In those years, developments failed to solve the problem of simultaneously achieving high strength and high workability to an alloy as well as providing safe long-term operation of designed protection coatings. This became a cause of certain stopping in carrying out researches in the field of Nb-alloys that was aggravated by successes achieved in those years in the field of designing superalloys and advanced cooling structures. In the middle of the 1980s, refractory metal alloys (especially niobium alloys) became an object of research interest again in connection with searching for materials characterised by high
Among the alloys based on refractory metals of groups VA and VIA, niobium alloys are perhaps the most varied. The niobium alloys used in various fields of engineering may be conventionally divided into two groups--structural and precision. To the first group, high-temperature heat-resistant alloys are attributed. To the second group, alloys having special physical properties are attributed. Among them are superconducting alloys; alloys characterised by specified values of linear thermal expansion coefficient, high elasticity and proper biocompatibility (for medicinal application as implants). The most numerous group of alloys to which application of niobium is oriented high temperature strength alloys. More than 30 years ago, in the 1960s, the attention of specialists engaged in material sciences was drawn to refractory metals as structural materials to be used in aerospace engineering. Apart from the ability to operate under high-temperature conditions, such materials should meet the following requirements: high workability and low density providing reduction of the launch weight. Concerning the above mentioned niobium with tmelt=2468°C, having the lowest density (8.4 g cm -3) among all refractory metals and the highest low-temperature ductility, the highest solubility of interstitial elements 303
E. N. Sheftel, O. A. Bannykh
304
Table 1. High-temperature niobium alloys and their strength levelsl,z Low-strength
B-33: Cb-33: C-103(KB-3): NTzU(KB-1):
5V 5V- 12.5Zr 10Hf- 1Ti-0.7Zr 1Zr- 0"IC o~, MVa
llO0°C 160-260 Medium-strength
B-66: FS-85: VN-3: 5VMTz: D-43:
1200°C .
0"100, M P a
.
1300°C
.
llO0°C .
High-strength
F-48:
F-50: VN-4: Series VAM WC 3009:
1300°C
1200°C 190-280
0"100, MPa
1300°C
llO0°C 160-240
1200°C 100-130
1300°C 56-70
15W- 5Mo- 1Zr- 0"1C 15W- 5Mo - 5Ti- 1Zr 10Mo - 1.5Zr - 0.3C - 0-03Ce,La (22 + 28)W- 2Hf-(0"067 + 0"13)C 30Hf- 9W a~, MPa
llO0°C 350-450
1200°C
5Mo - 5V- 1Zr 27Ta- 10W- 0.7Zr 4.6Mo - 1.4Zr- 0.12C 5W- 2Mo - 1Zr- (0"05 + 0.07)C 10W- 2"5Zr a~,MPa
1100°C 260-450
.
1200°C 180-350
Ol00,MPa 1300°C 140-300
llO0°C 210-300
1200°C 120-220
1300°C 100
*Annealed state. operating properties, to be used in new generations of spacecraft, aircraft and powder engineering, as well as in connection with development of new progressive technologies. In this search, a problem appeared in design materials intended for filling two operating temperature ranges. The first temperature range extends from 600 to 900°C where wrought Ti-alloys do not already work whereas Ni-alloys do not meet the specific strength requirements. The second one extends from 1100 to 1500°C where wrought Ni-alloys do not already work whereas alloys based on intermetallics and heat-resistant ceramics, considered as perspective structural materials, do not provide the required strength, ductility or fracture toughness. The present report, drawn up on the basis of further research carried out by the authors and the analysis of the published data, deals with the scientific principles of alloying and control of the structure of niobium alloys providing high strength and workability of the alloys for operating temperature ranges from 1100 to 1400°C and from 600 to 800°C. To design Nb-alloys destined to operate at temperatures where 0-1 Tin,it< T< (0"3 + 0"65) Tmelt all
the known methods of strengthening may be applied, such as the solid solution one, alloying by substitutional and interstitial elements, precipitation or dispersion strengthening, as well as creation of special structures. As a rule, the most effective combination of strength and ductility is achieved when joint application of several reasonably chosen methods of strengthening is used.
SOLID S O L U T I O N S T R E N G T H E N I N G Proceeding from the known mechanisms of solid solution strengthening, producing and some other requirements to be met by the designed alloys, the alloying elements should comply with one of the following criteria: (1) presence of wide solubility range even if at high temperatures; (2) an increase of intra-atomic bonds that is sufficiently reflected in increased melting point (Tme~t) of the alloy-strengthening in this case is kept at T> 0"5 Tmelt where the unstrengthening processes are determined by diffusivity of the elements; (3) a great difference of atomic sizes as compared to N b - this parameter determines the origin of considerable elastic distortions in the crystal lattice causing
Niobium--base alloys dislocation impeding the plastic flow process-such strengthening is kept at T< 0"5 Tmelt where the strengthening processes do not depend on diffusivity of alloying elements; (4) the alloy workability should not be radically decreased; (5) the alloy should not become considerably heavier (this is of great importance with respect to flying and rotating constructions) and should not be very expensive. W, Mo, Ta, and Re meet the requirements for operating temperature higher than 0"5 Tmelt.Content of these elements should not exceed 10-12 at.% to keep satisfactory workability of the alloys.3 A considerable difference of atomic sizes between Nb (Rat=l.46A) and Cr (1.256), V(1.31), Zr(1.59) and Hf(l'60) should provide strengthening of Nb with these alloying elements at operating temperatures below 0.5 Tme~t. Os, Ir, Ru and Re, meeting some of the requirements mentioned above, have high density (22.61, 22.65, 12.45 and 21.03 g cm -3, respectively) and are very expensive, therefore they should not be used for alloying Nb. Interstitial solid solution strengthening of niobium (by C, N, O, B) is not promising at high temperatures due to high diffusivity of these elements under these conditions. Moreover they decrease low-temperature ductility of alloys. These approaches based on analysis of binary systems are simplified to a certain extent. It is more reliable to estimate joint influence exerted by several alloying elements, i.e. on the basis of multicomponent systems.
STRENGTHENING BY DISPERSOIDS
The other method of increasing heat resistance of Nb is connected with precipitate and/or dispersion strengthening. Some oxide, carbide, nitride and boride dispersions as well as intermetallic ones or a combination of these phases can be used in this case. The strengthening phase should meet the following requirements: (1) higher mechanical strength as compared with that of the matrix; (2) particles must have sufficient size to be nonshearable obstacles to dislocation motion; (3) sufficient volume content to diminish Orowan loop; and (4) stability of particles size at operating temperatures. It is determined by high thermodynamic characteristics of the phases, its interaction with the matrix (components of the phase should exhibit low solubility and low diffusivity in the matrix) and low values of the surface energy with matrix.
305
The systems with two phase equilibria between Nb (or its solid solution) and such phases as intermetallics (Os, Ir, Ru, Pd, Pt, A1, Fe, Co, Ni and Cr), solid solutions based on alloying elements (Zr, Hf) or oxides, carbides, nitrides and borides both formed by Nb and the alloying elements can be revealed on the base of binary and ternary Nb phase diagrams. Intermetallic compounds are formed at relatively high concentrations of alloying elements (A1, Cr and Fe) radically lowering the temperature of Nb melting. So at temperatures exceeding 0.45-0-5 TrueStintermetallic phase particle coagulation is accelerated owing to an increase of diffusivity of the alloying element and self-diffusion of Nb. 4 Thus the precipitate strengthening by intermetallics can be effective for Nb below 1000°C.
STRENGTHENING BY OXIDES, CARBIDES AND NITRIDES
Strengthening of Nb at temperatures exceeding 1000°C should be provided by thermodynamically stable oxides, carbides, nitrides and borides. The most thermodynamically stable are oxides, carbides and nitrides of metals of Periodic System group IVA (Ti, Zr, Hf and Th) as well as the highest oxides of metals of group IIIA (A1, Y, La, Ce and Lu) and monoxides of metals of group IIA (Ba, Sr, Mg and Ca). When choosing phase, one should take into account compatibility of its elements with Nb minding an alloy producing process. So alloys with oxides of metals of group IIIA which are poorly soluble in Nb both in solid and liquid states, can be produced via powder technology. In this case the phases should keep thermal stability during compacting and sintering operations. Oxides, carbides, nitrides and borides of group IVA metals are dissociated and dissolved in the melt of high valency Nb ÷5 and during solidification are crystallized as separate phases or as solid solution of the elements in Nb. The alloys with these phases can be produced by melting. The researches, 5-15 accumulation and systematizationl 2,13,16-20 of phase diagrams Nb-Melv-X (X=O, N, C, B) showed that they are characterized by two types of isothermal section. The first (Fig. l(a) and (b)) shows that Nb can be in binary eutectic equilibria with MewX. It means that the phase diagram Nb-Mew-(O,N) has a pseudobinary eutectic section (Fig. l(d)). Presence of the ternary equilibria area Nb-Nb2X-Me~vX in
306
E. N. Shefiel, O. A. Bannykh
carbon-containing systems prevents the formation of a completely pseudobinary section (Fig. l(c)) but within a certain concentration range, along the Nb-MeivC direction, two-phase equilibria occur between Nb and (Nb,Melv)C. The second is a solid solution of NbC and MetvC. Such a Nb-MewC section may be considered as a partial pseudobinary one. Only those alloys which belong to the pseudobinary systems may exhibit the highest strength. Based on the laws of interaction between the phase diagrams and 'composition-high-temperature strength' diagrams 1,zl the highest strength belongs to those alloys of composition close to the solubility limit of MeivX in Nb (the I group of alloys) and to those which have hypoeutectic composition (the II group). The heterophase structure of the first of them is formed as a result of a solid solution decomposition whereas precipitation strengthening effect is determined mainly by dispersoid size. The heterophase structure of the II group of alloys is formed mainly during crystallization whereas strengthening is provided by a considerable growth of phase amount determined by the eutectic composition. 1 1 phase N ] 2phases Co)A ~3phases //" ~ )2 Nb~MeN
Nb
C
Meiv Nb
Melv
(d) ............... I
N Nb
Fig. 1.
i
btI~[~---II~ I I I ~ Groupsof alloys ,I MeW Nb
MeivX
Isothermic (a)-(c) and pseudo-binary (d) section of phase diagrams Nb-Melv-X (X = O,N,C).
The analysis of the own and published experimental data on concentration dependence of the strength of Nb-ZrO2(HfO2), Nb-ZrN(HfN) and Nb-ZrC(HfC) alloys and comparison of these data with the corresponding phase diagrams made it possible to designate the concentration ranges of alloys of the I and II groups (Table 2) in pseudobinary systems Nb-ZrOz(HfO2), Nb-ZrN (HfN) as well as in the system Nb-ZrC(HfC) the binary equilibria in which is realized not in the whole range of concentrations. Oxide strengthening
Oxide-containing alloys can be produced by charging the oxides during electric arc melting 16,17,22 as a result of 'contamination' during manufacturing and testing of material 23,24 by powder technology 25 or as a result of internal oxidization. 26-31 In ageing Nb-l% Zr(Hf)-O alloys the highest strengthening is achieved after quenching from 1700°C and following ageing at 950°C (Fig. 2). This effect is the result of precipitation of dispersed plate monocline phase ZrO2(HfO2) which is coherent to the matrix and oriented by its length a l o n g ( 1 0 0 ) N b . 16,17,22 In N b - T i - O 25 alloys the highest strengthening is achieved after ageing at 1000°C due to T i O 2 particles of 60 A size and with average distance between particles equal to 350A, coherent to matrix. 25 In the internal oxidized N b - l % Z r alloys while ageing at 1100°C, Z r O 2 particles of 40-60 A precipitate. 29-31 The thermal stability of MewO2 particle sizes in the ageing alloys depends greatly on what oxide forming element is present, in excess, in solid solution. Elevated content of Zr(Hf) in the solid solution above the stoichiometric M e i v O 2 facilitates the oxide particles coagulation. The accelerated growth of particles at 1100°C in the alloys with excess Z r 16,17,24 (Table 3) arises from the presence of free Zr(Hf) controlling the speed of growth of the phase particles and from
Table 2. Content of the dispersoids providing the precipitation (I group) and the dispersion (II group) strengthening in Nb-Meiv-X (X = O,N,C) alloys
System
Nb-ZrOz(HfO2) Nb-ZrN(HfN) Nb-ZrC(HfC)
Optimal content of dispersoids in alloys of I group
Content of dispersoids in alloys of H group
(rot. %)
(vol. %)
0.5-0"6 for at.%Zr/at.%O > 0"5 2 for at.%Zr/at.%O -<0-5 5 3"5
1-2-2 5-15 3-17
307
Niobium--base alloys /•..._.-e _•/
,, ~
•f ° /
Table 4. The size of ZrO2 particles in internal oxidized N b - l % Z r alloy annealed at 1600°C versus the free oxygen concentration in solid solution 29-31
Tq = 1700°C t = 2h --
\
o .~ y_.
.,.-~
-2 ~ .~
Annealing time
Diameter of particles (A)
~h)
Free oxygen concentration (at. '%)
-4 2800"
1
2400"
10 30
Z~
0"16
0"5
1.7
100 170 230
100 160 190
90 120 160
> ~Z 2100 1200 16_18x ~ 1700
1300! 600
_~ 1000 ~ ~'~, ~ _
800 10130 1200 Temperature of ageing (°C)
~8 0 06001~
~
o [] • • A ©
2% ZrO2,int. oxid. 1.5% ZrO2, aged 1% ZrO2, aged 0.3% ZrO2, recryst. 0.3% ZrO2, aged 0.12% ZrO2, aged
~ 6 " - 6 ,
%
Fig. 2.
Effect of quenching and following ageing temperatures on hardness and variation of resistivity for two-phase N b - 1 wt.%ZrO2 alloys with 0-3 vol.% Z r O 2.
"2
20x 400 ~•
Table 3. The size of Z r O 2 particles in N b - l % Z r - O alloyst with (0"3 + 0.6 ) vol.% Z r O 2 versus the ageing conditions
Ageing Plate length, Plate thickness conditions (A) (A)
Cc)
cA)
950
5
30
950
20
50
950 100
60
1100 1200
Distance between particles
20 5
150 570
Several atom layers Several atom layers Several atom layers 20 30-50
--
\ 18x
200
3 29 x
"A.~ I I 800 1000 Test temperature (°C)
I 1200
Fig. 3. UTS and elongation for two-phase N b - l w t . % Z r - O alloys with various contents of ZrO2 versus test temperature.
-70-140 200 400-860
t A f t e r quenching from 1700°C.
the increase of oxygen diffusivity in Nb at 1100°C. This results in reducing the effect of precipitation strengthening at this temperature. An efficient suppression of oxide particle coagulation up to 1600°C has been obtained in internal oxidized alloys with excess oxygen content (related to MetvO2)in solid solution (Table 4). The authors 29-3~ explain this fact by dramatic decrease of equilibrium solubility of ZrO2 in the solid solution in presence of excess oxygen. In order to provide sufficient ductility of alloys free oxygen content should not exceed 0.2-0.3 at.%. The structure of these internal oxidized alloys is characterized by presence of 1 vol.% of highly dispersed (40-60 A) coherent to the matrix particles of f.c.c, modification ZrO2(HfO2) precipitated from the solid solution. High elasticity
modules and high dispersivity of this phase preserved up to high temperatures, provides a considerably greater increase of recrystallization temperature in the internal oxidized alloys (up to 1 8 0 0 ° C ) 26'27 a s compared wth alloys having similar amounts of phase but with excess free Zr in solid solution (122 5°C) ~7and a greater effect of precipitate strengthening (Fig. 3). Oxide strengthening increases the creeprupture strength and creep resistance of alloys16,17,23,24.27,28,32,33 greater than tensile strength (Figs 3 and 4). The creep resistance is extraordinarily sensitive to extremely low oxide dispersoids content.23,24,2s,34 Oxide strengthening alloys exhibit satisfactory ductility (~p,6>_10%, Tbr_<-50°C)within the temperature range 20-1300°C (Fig. 3). This can be explained by a low volume content of oxide particles ( < 1 vol.%), their high dispersivity and small grain size. Exceeding 1.5 vol.% of oxide phase in alloys leads to formation of large (2-3 pm) hypereutectic oxide particles and to a
308
E. N. Sheftel, O. A. Bannykh
10° I
Table 5. The size of particles of alloys with ZrN versus ageing conditions
o IO00*C = 70MPa ol
--,
•
Time of ageing (h)
2
10-1
Size of particles (.4) Temperature of ageing (*C) 900
1-3 20 (") 10"2
30-60
0
1200
1400
1500
30-90 120-140 300-500 200-350 70-120 200-270 150t 80-200 250-330 120-300 600-800
50 1-3
tDistance between particles.
10-3 0
1000
I 0.3
I 1.0 Vol%Z r O 2
I 1.5
Fig. 4. Creep rate of two-phase Nb-lwt.%Zr-O alloys versus volume content of ZrO2: 1--recrystallize annealing; 2--quenching 1700"C + ageing 950"C, 10 h.
7(
great loss of strength. In this case the alloys are greatly embrittled. Nitride s t r e n g t h e n i n g
o
5(
r~
Nitride containing alloys can be produced by arc melting of nitride containing charge 16,17,35 or by melting in nitrogen containing plasma, 36 as well as by the diffusion saturation method 37 and powder process. Investigations 38-4° of ageing in Nb-ZrN(HfN) alloys have showed that active solid solution decomposition occurs within the temperature range 900-1200°C and leads to formation of 5-6vo1.% of highly dispersed f.c.c. ZrN(HfN) particles* (Table 2). The greatest effect of strengthening is achieved after quenching from 1700°C and following ageing at 1000°C. The precipitates of ZrN(HtN) after this condition in this case are of 150-450 A, the distribution density is (5-7)× 1016 particles cm -3, coherent with the matrix. Z r N ( H ~ ) particles stay highly dispersed up to 1500°C (Table 5). Thermal stability of a two-phase structure at 1000-1500°C exerts influence upon development of recrystallization processes occurring in the deformed alloys and taking place within the same temperature range. Highly dispersed nitrides (Table 5) slow down transformation of deformed :~Owingto formation of solid solutions between MewN and NbN in Nb-Meiv-N systems (Fig. l(a)) nitride (MewNb)N should be formed, however, presence of Nb in the nitride has not been determined experimentally probably because of its low content.
31
800
1000 1200 Test temperature(*C)
Fig. 5. UTS divided by alloy density for some commecial Nb-base alloys and for two-phase Nb-Zr-N alloys with various volume contents of phase ZrN versus test temperature.
structure and completion of the primary recrystallization up to 1400-1500°C. The nitride phase effectively slows down grain growth. The grain size in the alloys with 5 vol.% nitride is less than 15-20/~m after annealing at 1700-18000C for 1 h. 41 Tensile strength and rupture strength of the nitride containing alloys is increased within the temperature range 20-1200°C with an increase of the phase amount and its dispersivity 12,13,3s,39,42 and depends on heat treatment conditions.12,13,16,17, 37,42 Specified heat treatment gives these alloys strength and specific strength larger than that of many high-strength Nb-alloys with highly alloyed solid solution (Fig. 5, Table 6).
Niobium--base alloys It is known that the phase composition of the alloys obtained under real conditions may not comply with the stable equilibria phase diagram. This may be connected with stability of the metastable state or with deviation from the specified chemical composition. It is by these reasons that one can explain occurrence of large eutectic nitrides up to 3000/k in the structure of precipitation strengthened alloys. In these alloys, besides ZrN(HfN), hexagonal phase N b 2 N 3 8 - 4 ° having plate shape with average linear dimensions of (2000 + 10 000)/kx (300 + 1200) ,/k is released at the earlier stages of ageing at 800 to 1200°C. Nitride Nb2N is not stable and it dissolves in 2-5 h after its precipitation. Neither eutectic nitride or Nb2 N contributes to strengthening at 900-1200°C. Information about the alloys of the II group is limited by the alloy with 12vol.%ZrN. 14 The eutectic type structure of this alloy is represented by ZrN fibres 0-8-1 #m in thickness and up to 20 #m in length. Tensile strength and rupture
Table 6. R u p t u r e stress (o~00) of Nb-alloys with nitride strengthening
Alloy
Oloo, MPa Annealing at 1400°C
strength of this alloy is lower than that of the alloy with 4 vol.% ZrN up to 1300°C.
Carbide strengthening A great number of Nb-alloys developed in the 1960s and 1970s are based on precipitation strengthening by carbide (Table 1). The fact that at a Zr content within 1-2 mass% and C, 0.12-0.13 mass% (-1-1-5vo1.% of carbide phase) alloys characterized by satisfactory workability have predetermined compositions of the alloys developed in that time (Table 1). Formation of two-phase structure in the alloys with 1-2% carbide during cooling from the singlephase state proceeds with precipitation of the hexagonal carbide N b z C ( o r with Nb3C 2 at rapid cooling) in form of plates of 0.5-1 #m size. 15,43-45 It is not possible to suppress precipitation of these carbides even at such rates of cooling as ( 5 - 7 ) X 106 K s - l . 46 Nb2C is not stable at high temperature, while being heated up to 1000-1200°C it is transformed into stable f.c.c. carbide containing Mew, with a lattice parameter close to that of NbC (4-47 A). Completeness of transformation Nb2C to f.c.c, carbide as well as Zr or Hf content in it depends on ageing temperature a n d d u r a t i o n . 43,47
Quenching from 1700°C + ageing at IO00°C, 20 h
Test temperature (°C)
Nb- lvol.%ZrN Nb-2-5vol.%ZrN Nb-4vol.%ZrN Nb- 10%Mo-4vol.%ZrNt N b - 1-5vol.%HfN Nb-4"5vol.%HfN Nb-I 0%Mo-4vol.%HfNt
1100
1200
1100
1200
-120 ------
-65 ------
160 240 255 -----
90 110 130 140 100 145 185
? Q u e n c h i n g f r o m 1800°C + ageing at 1000°C, 20 h.
Phase transformations occurring in alloy with 1 mass% Zr and 0.12 mass% C were investigated firstly in Refs 43 and 47. Developed heat treatment conditions have provided increasing the strength of such alloys in ever higher degrees (Table 7). The alloys with 1-1.5vo1.% carbide keep their strength up to 1100-1200 ° (Fig. 6). Information concerning alloys with carbide phase content higher than 1-2 vol.% is limited by the data concerning alloys VN-4; 1 P W C - 3 3 49 a s well as by the results of analyses made by the authors j0-52 Arc melted alloys containing up to 12 vol.% carbide dispersoids have been investigated in Refs
Table 7. M e c h a n i c a l p r o p e r t i e s o f N b - 1 0 % W - 5 % M o - l % Z r - 0 - 1 2 % C
Treatment
Forging Quenching f r o m 1700°C Quenching + ageing LTMTt
309
alloy after various t r e a t m e n t s (by O. B. Dashevskaya)
Mechanical properties at 1100°C oB (MPa)
ao.2 (MPa)
6 (%)
-46
43 28
-31
44
29
35
68
56
17
t L o w - t e m p e r a t u r e thermo-mechanical treatment.
Oloo (MPa) 24 --
28.5
Life-time (h) 60 1200
310
E. N. She#el, O. A. Bannykh 500 Nb - 15W - 5Mo - 1Zr - 0,1C (F-48) 450 -
%-
~4ooe~
350 -
300 o
E .=
250 -
200-
150 1050
I
I
I
I
I
1100
1150
1200
1250
1300
Test temperature (°C)
140
¢•-
Fig. 7. Microstructure of the Nb-7vol.%(Nb,Zr)C alloy. Optic micrographs: (a) as cast Vcoot=10 ~ K s-l; (b) after hydrostatic extrusion with counter-pressure, 25% reduction; (c) see (b), 60% reduction; (d) ribbon of 150 # m thick, Vcool= 105 K s-l; (e) transmission electron micrograph, see
10W - 1Zr - 0.1C (D-43)
120
1~--
(d).
80--
Table 8. Size of carbides (Nb,Zr)C particles in Nb alloy with 6 vol.% carbide versus annealing conditions
Nb - 10W - 0.15C
20 1050
I
I
I
I
I
1100
1150
1200
1250
1300
Test temperature (°C)
Fig. 6.
Annealing conditions
(°c)
Mode of particles
(h) Eutectic,
(#m)
Result of dispersion strengthening of Nb alloys by 1.2 vol.% carbide phase (Nb,Zr)C.
50-52. The phase composition of the alloys is represented by Nb-base solid solution and f.c.c. carbide (Nb,Zr)C being in eutectic equilibria. Their as-cast structure is represented by large, 5 - 1 0 # m , nonstoichiometric eutectic carbides (Nb,Zr)xC1-x with a lattice parameter close to that of NbC (4.47 A). Zr concentration in the carbide and a degree of its approximation to stoichiometric composition increase with increasing Zr and C content in the alloy. Being precipitated primarily at grain boundaries, these carbides form an almost continuous network along the boundaries (Fig. 7(a)). Annealing of cast alloys leads solid solution decomposition and precipitation of (Nb,Zr) C particles oriented by their length along {100} Nb (Table 8). Alloys with 5-12vo1.% carbide have much higher strength and specific strength up to 1400°C in comparison with those of alloys containing 1-2% of the phase (Fig. 8). They have high
Particle size
1000 1200 1200 1200 1200 1600 1800 2100 2100
1 1 5 20 50 1 1 1 1
2-8t 2-8t 2-8 2-8 2-8 2-8 2-8 up to 6 up to 6
Segregatedfrom solid solution Plate length (,4)
Platewidth (.4)
100-150 600-800 1000-1500 1300-1700 2300-3400 1900-2500 5000 ---
20-30 50-100 200-300 200-400 350-400 300-400 500-600 ---
tit depends on crystallization rate.
capacity for strain hardening and high recrystallization temperature (1500°C). However, a disadvantage of these alloys is their poor ductility and fracture toughness at room and medium temperatures, which prevents their application. One of the main reasons for it is large carbide particles making up a continuous network along grain boundaries (Fig. 7(a)). During plastic deformation local stress concentration at large carbide particles promotes local fracture away from the matrix (which occurs more easily at matrix-particle
Niobium--base alloys × ",~ • I zx 120"" • " [ - xD
Mo, 5vo1% HfN, 80/~ 1.5% ZrO2, int. ox., 100/~ Mo 4% ZrN, 80/~ 4%'ZrN, 80/~ 6% (Nb Zr) C h-extr. (2000 x 300) ,~ Mo, 6%'(Nb, Zr) C, cast (4000 x 400) .~
1000 --o x~
g ~o~ 52" .-~.."-~,L--a--..×sx
:..;&5,,,\%on
600
- - ..--
-...
o~ o
NX×
11 x
- .,2,X~xl2X ~ ~i5 x
400
if-
~ 2000
- " " ~
r~37x
~ ~ " ~ " - - -'~e~N• 1.4% Zr02, 5500A " ~ ' ~ ' ~ - ~ . o 0.3% ZrO2, 40A • O 1,~% (Nb, ~r)C, (1300 x 250)i~ 700
Fig. 8.
~.
Comparison
900 1100 1300 Test temperature (°C)
1500
o f oxide, n i t r i d e a n d c a r b i d e s t r e n g t h e n i n g effect.
interface). As a result, the cracks nucleate at these sites. 5° Increase of low-temperature ductility of cast alloys with high carbide content can be achieved via optimization of cast structure by means of proper heat treatment, high-temperature extrusion (1700-1800°C) or by applying such processes as cold hydroextrusion with counterpressure ~7.50and rapid solidification.5°-52 Multitransitional hydroextrusion with counterpressure (30-15 kbar) made it possible to deform the alloys with 5-7 vol.% of the phase with reduction of up to 80-90%. The deformation leads to breaking of large eutectic carbides, distintegration of the grain boundary continuous carbide network and to increasing of structure uniformity (Fig. 7(b) and (c)). The hydroextrusion leads to the formation of a preferred texture of (110) type in the extrusion direction. The above factors improve material ductility. Along with the increase of ductility, hydroextrusion results in a considerable strain hardening that is retained up to 1500°C. Rapid solidification of the alloys with 5-9vo1.% carbide, carried out by different methods providing a cooling rate of 105 K s -1, leads to refining of the structure (Fig. 7(d) and (e)). The grain size after rapid solidifcation is 1-3/~m, and that of eutectic carbides--0.2-0.3 ~m at decreasing of their total amount and disappearance of grain boundary continuous carbide network (Fig. 7(d) and (e)).
311
All the above described should provide an increase of strength (according to Hall-Petch) and ductility. Increasing ductility is provided by facilitating propagation of slip from one grain to an adjacent one across the 'feed' boundary and decreasing the contribution from matrix-particle interfaces along which the nucleation of cracks occur. Annealing of rapidly solidified alloys at temperatures up to 1800°C does not lead to grain growth and coagulation of eutectic carbides but leads to decomposition of the solid solution with formation of the dispersed (Nb,Zr) C (Table 8). The above results allow us to suppose that the material produced via rapid solidification followed by compacting would exhibit a unique combination of strength and ductility. The alloys of eutectic composition produced by directed crystallization deserve attention. Directionally crystallized Nb-Nb2C alloy containing up to 30vo1.% carbide as fibres have the specific strength of 4.5 km at 1200°C. 53 Directionally crystallized Nb-(Nb,Mew)C eutectic alloy possesses the specific strength of 6 km at 1200°C exceeding that of all known Nb-alloys. Comparison of oxide, nitride and carbide strengthening Comparison of the above, within 20-1400°C, has been made taking into account superposition of such factors as thermodynamic stability, quantity and dispersivity of the strengthening phases (Fig. 8). Thermodynamic stability of the phases increases in the order: carbides, nitrides, oxides. The free energies (AH~98,kJ mol-1) a r e : N b C 140"3; Z r C - 200; H f C - 230.1; Z r N - 365"3; HfN - 369"2; ZrO 2 - 1094.1; HfO 2 - 1113. The amount of strengthening phase is increased from oxides to carbides and then to nitrides for I group alloys and from oxides to nitrides and carbides for II group alloys (Table 2). In I group alloys the highest strengthening up to 1200-1300°C is provided by nitrides formed in the alloys after quenching followed by ageing in the greatest amount (4-5 vol.%) and the highly dispersed state (70-90 A). The alloys with 1-2.5vo1.% carbides exhibit the lowest effect of strengthening. The cause of this is the low thermodynamic stability of the phase formed in these alloys during ageing (NbC) and the relatively large size of its particles. Strength of precipitate strengthened alloys with -< 1.5 vol.% ZrO2(HfO2) is intermediate between those of carbide and nitride containing alloys. But
312
E. N. Sheftel, O. A. Bannykh
the strength can be increased up to that of nitride containing alloy by providing small size ( < 100 A) oxide particles up to 1600°C due to an excess content of oxygen (related to ZrO2(HfO2)) in solid solution (Fig. 8). The strength of alloys with 1-2vo1.% (Nb,Melv)C is inferior to that of alloys with oxides and nitrides, but slower with temperature increasing degradates. The latter is provided by lower solubility of carbides in Nb matrix and slower coagulation of multicomponent phase (Nb,Mew)C particles by comparison to two-component oxide or nitride particles. The alloys with 5-9 vol.% carbide dispersoids exhibit greater high-temperature strength than alloys with 1-2 vol.% carbide, but are inferior to precipitate strengthened alloys with nitride. The strengthening by phases with Hf is more effective than that by phases with Zr. JOINT SOLID SOLUTION AND OXIDE, NITRIDE AND CARBIDE STRENGTHENING Alloying of Nb by W, Mo, Ta and V, as well as by Melv and one of the interstitial elements (O,N,C) provides joint solid solution and dispersoid strengthening at 20 to 1400°C (Fig. 9). The contribution from each of them to the general strength level depends on alloy chemical composition, phase content and testing (operating) temperature. Thus, at 1300-1400°C (0-55-0.6 Tmelt) when solid solution strengthening in Nb-alloys is being decreased, the contribution of dispersion strengthening by carbide and nitride becomes important.
900[ - - ~ 800~ ~
~= 700l--
UTS attainable by MelvC,MeivN ~
/
~ 600q) 500'x'- ~
Two-phas,~ N
* 400 --~1 Single-phase~ | "~ ~~ ,.-- ~ alloys 300-I"
__,~, 600
Fig. 9.
I
~/~///A~ W(,ff//~7/fl~
I
1000 1200 Test temperature(*C)
1400
Ultimate tensile strength of single-phase and twophase Nb-alloys.
Solid solution alloying by W of Nb-MewX alloys leads to a decrease of their low-temperature plasticity, 12,13 while alloying by Mo leads to an increase. 54 Simultaneous increase of tensile strength, stress-rupture strength and plasticity of alloys with Mo (Fig. 8) are provided by solid solution strengthening and Mo influence on the twophase structure formed in the alloys. The influence has been carried out by following the refining of grain size, decrease of interstitial element content in solid solution providing increase of precipitating phase amount as well as increase of strength of particle-matrix interface. PRECIPITATON STRENGTHENING BY INTERMETALLIC PHASES Increasing demands for new materials used in aircraft engine construction, conditioned by the elevation of operating temperatures of the engine, stimulate a search for the materials which can fill the operating temperature gap from 600 to 800°C. At these temperatures, high temperature strength Ti-alloys can not work whereas Ni-alloys do not meet the specific strength requirements. Russian scientists have developed high-temperature, heat-resistant, wrought alloy, operable at 600-700°C. The principle of its alloying lies in the choice of elements providing formation of protective films which increase resistance of the alloys against oxidization (Ti,AI) and formation of precipitate strengthening phase (Ti3A1). The specific yield strength of the alloy is considerably higher than that of Ni-high-temperature wrought alloy operable at 600-800°C. When stress-rupture tested in air, the alloy exhibits significant stress riser (notch) sensitivity at 600-700°C. Moreover, specific stress-rupture strength of the alloy becomes considerably lower than that of Nialloys. The cause of this phenomenon is the oxidizing effect of the environment. 55 The notch sensitivity was shown to be controlled by oxygen that is absorbed from the environment to form a special underscale layer of altered chemistry and phase constitution. The environmental damage could be prevented to a large extent by optimization of the structure and alloying with carbon. A perspective of heat-resistance increasing of high-temperature alloys of the Nb-Ti-Al system may be their additional alloying by Cr. The dispersed intermetallic phase NbCr 2 is expected to be more stable against coagulation in Nb-matrix than Ti3AI and therefore to provide more effec-
Niobium--base alloys 1000 _-800 v
(VM4ArrAN)
600 --
~'5oo
_
~
\
REFERENCES
Mo-alloys (TzM-5, TzM-6)
~
~
v ~ L N - 1 TMT . ~_ 5VMTz
] 1
LN - ~
!
~
~ i n t . oxid. f ~-(~o, rio C
directed eutectic
/ e 200- (IN-100)-/~x ~
"~////_/(/
Ni-~loys, directed ~ V ~ eutectic ~ ~ .
~
\\'v~(g,f~//~
_
,
._ \ _ X
-
-
4
".
Nb-9Mo-4ZrN D-43
o ---~, 700
I
900
I
I
1000 1100 Test temperature (°C)
I
1200
I
1300
Fig. 10. The 100 h stress rupture strength of Nb-base alloys in comparison with high-strength Ni- and Mo-base alloys.
tive strengthening. Cr is able to increase the heatresistance of Nb-Ti-AI alloys by 100°C to 800°C, due not only to Cr-containing protective surface film formation, but also to holding of Nb in NbCr2, decreasing the amount of Nb moving towards the surface layer.56
CONCLUSION The above described data and concepts can be the basis for the most efficient choice of Nb-base alloy compositions, methods and conditions of production and treatment, providing high-temperature strength, workability and, for medium operating temperatures, high heat-resistance. This approach allows development of Nballoys with high-temperature strength, higher than that of Ni-superalloys above 800°C and similar or even higher than that of Mo-based alloys at 1000-1300°C (Fig. 10). This group of Nb-alloys requires special coating protecting against oxidization. The approach also provides for development of Nb-alloys to be used at medium temperature, being heat resistant, with a specific strength at 600-800°C higher than that of high-temperature wrought Ni-alloys.
1. Zakharov, M. V. & Zakharov, A. M., In High Temperature Strength Alloys. Metallurgija, Moscow, 1972, pp. 331-5. 2. Loria, E. A., J. Metals, 7 (1987) 22-6. 3. Vasiljeva, E. V. & Prokoshkin, D. A., Metallovedenie i Termicheskaja Obrabotka Metallov, 8 (1967) 2-9. 4. Povarova, K. B., In Heat-Resistant and Oxidation Resistant Metallic Materials. Nauka, Moscow, 1987, pp. 56-76. 5. Ordanjan, S. S., Avgustinnik, A. I. & Vigdorovich, V. S., Zhurnal Prikladnoj Khimii, 39(2) (1966) 312. 6. Taylor, A. & Doyle, N. I., J. Less-Common Met., 13(5) (1967) 511. 7. Taylor, A. & Doyle, N. I., J. Less-Common Met., 13(4) (1967) 413. 8. Barabash, O. M. & Shurin, A. K., Izv. Acad. Nauk SSSR. Met., 5 (1978) 195. 9. Ono, K. & Morijama, J., J. Less-Common Met., 79(2) (1967) 255-60. 10. Shurin, A. K., Dmitrieva, G. P. & Panarin, V. E., Metallofizika, 6 (1979) 74-7. 11. Shurin, A. K. & Loktionova, V. A., lzv. Acad. Nauk SSSR Met., 1 (1970)231-3. 12. Trefilov, V. I. & Moiseev, V. F., In Dispersion Particles in Refractory Metals. Naukova Dumka, Kiev, 1978, pp. 122-35. 13. Trefilov, V. I. & Moiseev, V. F., In Dispersion Particles in Refractory Metals. Naukova Dumka, Kiev, 1978, pp. 148-9. 14. Shurin, A. K., Kozinskij, G. J., Sulzhenko, V. K. & Barabash, O. M., Doklady Acad. Nauk SSSR, 222(6) (1975) 1332-4. 15. Druzhinin, I. I., Zakharov, A. M. & Novikov, I. I., Izvestija Vuzov. Tsvetnaja Metallurgija, 1 (1972) 118-21. 16. Grigorovich, V. K. & Sheftel, E. N., In Dispersion Strengthening of Refractory Metals. Nauka, Moscow, 1980, pp. 138-70. 17. Grigorovich, V. K. & Sheftel, E. N., In Dispersion Strengthening of Refractory Metals. Nauka, Moscow, 1980, pp. 175-276. 18. Shurin, A. K., Poroshkovaya Metallurgija, 1 (1973) 50-3. 19. Shurin, A. K. & Dmitrieva, G. P., Metallofizika, 53 (1974) 91-5. 20. Shurin, A. K. & Panarin, V. E., Metallofizika, 66 (1976) 85-91. 21. Grigorovich, V. K., In High-Temperature Strength and Phase Diagrams. Metallurgija, Moscow, 1969, pp. 185-213. 22. Sheftel, E. N. & Grigorovich, V. K., In Alloys of Refractory and Rare Metals, ed. E. M. Savitsky & K. B. Povarova. Nauka, Moscow, 1984, pp. 127-35. 23. Pollock, W. J., Met. Eng. Quart., 2(4)(1962) 58-68. 24. Hobson, D. O., Metallurg. Soc. AIME, 18 (1963) 224-325. 25. Gennari, V., In Proc. 1st Europ. Conf. Adv. Mater Processes, ed. H. E. Exner & V. Schumacher, Vol. 2, DGM lnformationsgesellshaft, Verlag, 1990, pp. 731-6. 26. Dushin, Yu.A., Emel'janova, L. I., Novilov, L. D., Izv. Acad. Nauk SSSR Met., 5 (1978) 204-9. 27. Vedenjapin, G. A., Dushin, Yu.A., Emel'janova, L. I., Novikov, L. F. & Parhuta, G. N., Fizika Metallov i Metallovedenie., 47 (5) (1979) 972-7. 28. Tsviluk, I. S., Pilnokov, V. I., Avramenko, D. S. & Konovaitsev, A. V., Problemi Prochnosti., 9 (1984) 49-53.
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29. Korotaev, A. D., Tumentsev, A. N. & Sukhovarov, V. F., In Dispersion Strengthening of Refractory Metals. Nauka, Moscow, 1989, pp. 95-104. 30. Korotaev, A. D., Tumentsev, A. N. & Sukhovarov, V. F., In Dispersion Strengthening of Refractory Metals. Nauka, Moscow, 1989, pp. 121-8. 31. Korotnev, A. D., Tumentsev, A. N. & Sukhovarov, V. F., In Dispersion Strengthening of Refractory Metals. Nauka, Moscow, 1989, pp. 144-51. 32. Bonesteel, R. M., Roweliffe, D. I. & Tietz, T. E., Trans. Jap. Inst. Met., 9 (1968) 597-602. 33. McCoy, H. E., J. Less-Common Met., 8(1) (1965) 20-35. 34. Tsviluk, 1. S., Pilnokov, V. I. & Men'shikh, V. A., Problemi Prochnosti., 11 ( 1978) 65 - 8. 35. Anton, D. L., Snow, D. B. & Giamei, A. E, In High Temperature Niobium Alloys, ed. John Stephens & Iqbal Ahmad. TMS, Warrandale, Pennsylvania, 1991, pp. 49-72. 36. Erokhin, A. A., Pankratov, N. N., Pugin, A. I., Moiseev, V. F. & Ban'kovskij, O. I., Fizika i Khimifa Obrabotki Materialov, 1 (1973) 15-20. 37. Lakhtin, U. I., Kogan, Ja.D., Bobok, A. N., Zakharova, T. V. & Abrasin, A. G., Metallovedenie i Termicheskafa Obrabotka Metallov, 10 (1973) 13-17. 38. Sheftel, E. N., Usmanova, G.Sh. & Grigorovich, V. K., Izv. Acad. Nauk SSSR Met., 5 (1980) 165-9. 39. Sheftel, E. N., Grigorovich, V. K. & Usmanova, G.Sh., In Research and Use of Refractory Metal Alloys, ed. E. M. Savitskij, M. A. Tylkina & N. K. Kazanskaja. Nauka, Moscow, 1983, pp. 119-30. 40. Sheftel, E. N., Grigorovich, V. K. & Usmanova, G.Sh, Izv. Acad. Nauk SSSR Met., 1 (1975) 170-8. 41. Sheftel, E. N., Usmanova, G.Sh. & Grigorovich, V. K., Metallovedenie i Termicheskafa Obrabotka Metallov, 7 (1979)23-7. 42. Sheftel, E. N., Usmanova, G.Sh. & Grigorovich, V. K. & Geminov, V. N., Metallovedenie i Termicheskafa Obrabotka Metallov, 2 (1984) 48-51.
43. Bollenrat, E & Ostermann, E, In High Temperature Materials, ed. E Benesovsky. Springer-Verlag, Wien, 1969, pp. 130-50. 44. Sheftel, E. N., Grigorovich, V. K. & Trankovskaja, G. R., Izv. Acad. Nauk SSSR Met., 1 (1973) 205-11. 45. Ostermann, E, J. Less-Common Met., 25 (1971) 243-56. 46. Savitsky, E. M., Ivanova, K. N. & Revjakin, A. V., Doklady Acad. Nauk SSSR, 213(6)(1973) 1289-91. 47. Savitsky, E. M. & Dashevskaja, O. B., lzv. Acad. Nauk SSSR Met., 3 (1967) 152-8. 48. Technical Recommendation on Thermo-Mechanical Treatment of Niobium Alloy. Specification 14-127-8276. 49. Delgrosso, E. G. & Kaminsy, J. J., Pat. USA R75/175, (1965) S.N 327552, No. 3282690. 50. Bannykh, O. A. & Sheftel, E. N., In High Temperature Niobium Alloys, ed. John Stephens & Iqbal Ahmad. TMS, Warrendale, Pennsylvania, 1991, pp. 73-82. 51. Sheftel, E. N., Bannykh, O. A., Usmanova, G.Sh. & Markova, E. V., Metal. i Termich. Obrab. Met., 4 (1989) 32-5. 52. Sheftel, E. N., Bannykh, O. A., Usmanova, G.Sh., Usikov, M. E & Sujazov, A. V., Izv. Acad. Nauk SSSR Met., 6 (1989) 85-9. 53. Zalkind, A. E, Lemke, E & George, F., In Monocrystalline Fibres and Materials Reinforced with Them, trans. ed. A. T. Tumanov, Mir, Moscow, 1973, pp. 332-78. 54. Sheftel, E. N., Grigorovich, V. K., Usmanova, G.Sh., Liberov, Yu.E & Barabash, O. M., Izv. Acad. Nauk SSSR Met., 4 (1983) 121-7. 55. Sheftel, E. N., Bannykh, O. A., Liberov, Yu.E, Filip'eva, O. A. & Yudkovsky, S. Y., In Proc. 13th Plansee Seminar, Metallwerk Plansee, Reutte, ed. H. Bildstein & R. Eck, Vol. 1, 1993, pp. 757-71. 56. Bannykh, O. A., Sheftel, E. N., Usmanova, G.Sh., Sharapov, A. A., Kaputkin, D. E., Metally, in press.
Niobium--Base Alloys E. N. Sheftel & O. A. Bannykh Baikov Institute of Metallurgy, Russian Academy of Sciences, Moscow, Russia (Received 25 March 1994; accepted 16 May 1994) Abstract: On the basis of the research carried out by the authors and the analysis of the published data, the scientific principles of the alloying and designing of structure of niobium alloys, providing high strength and workability of the alloys intended for two operating temperatures:from 1100 up to 1400°C and from 600 to 800°C, have been considered. The developing principles are based on the revealing of interdependence between phase diagrams, thermodynamiccharacteristicsof elementsand phases formingin the proper alloys as well as their structures and mechanical properties. Main attention has been paid to the alloysof Nb-Mew-X (X = O, N, C) systemswith oxide, nitride and carbide dispersoids.
INTRODUCTION
and, connected with it, the least sensibility of ductile properties to their content, it was considered as the most perspective basis for structural alloys to be used under high-temperature conditions. At that period the principal problems for Nballoys to be solved have been in searching for ways to increase high-temperature strength at satisfactory workability, as well as methods of protection against oxidisation at operating temperatures exceeding 1000°C. As a result, at the end of the 1960s and beginning of the 1970s, a great number of alloy compositions were developed which may be conventionally divided into low-, medium- and high-strength alloys that have been used for a long time operating at 1000-1200°C and a short time working at 1200-1700°C (Table 1). t,2 In those years, developments failed to solve the problem of simultaneously achieving high strength and high workability to an alloy as well as providing safe long-term operation of designed protection coatings. This became a cause of certain stopping in carrying out researches in the field of Nb-alloys that was aggravated by successes achieved in those years in the field of designing superalloys and advanced cooling structures. In the middle of the 1980s, refractory metal alloys (especially niobium alloys) became an object of research interest again in connection with searching for materials characterised by high
Among the alloys based on refractory metals of groups VA and VIA, niobium alloys are perhaps the most varied. The niobium alloys used in various fields of engineering may be conventionally divided into two groups--structural and precision. To the first group, high-temperature heat-resistant alloys are attributed. To the second group, alloys having special physical properties are attributed. Among them are superconducting alloys; alloys characterised by specified values of linear thermal expansion coefficient, high elasticity and proper biocompatibility (for medicinal application as implants). The most numerous group of alloys to which application of niobium is oriented high temperature strength alloys. More than 30 years ago, in the 1960s, the attention of specialists engaged in material sciences was drawn to refractory metals as structural materials to be used in aerospace engineering. Apart from the ability to operate under high-temperature conditions, such materials should meet the following requirements: high workability and low density providing reduction of the launch weight. Concerning the above mentioned niobium with tmelt=2468°C, having the lowest density (8.4 g cm -3) among all refractory metals and the highest low-temperature ductility, the highest solubility of interstitial elements 303
E. N. Sheftel, O. A. Bannykh
304
Table 1. High-temperature niobium alloys and their strength levelsl,z Low-strength
B-33: Cb-33: C-103(KB-3): NTzU(KB-1):
5V 5V- 12.5Zr 10Hf- 1Ti-0.7Zr 1Zr- 0"IC o~, MVa
llO0°C 160-260 Medium-strength
B-66: FS-85: VN-3: 5VMTz: D-43:
1200°C .
0"100, M P a
.
1300°C
.
llO0°C .
High-strength
F-48:
F-50: VN-4: Series VAM WC 3009:
1300°C
1200°C 190-280
0"100, MPa
1300°C
llO0°C 160-240
1200°C 100-130
1300°C 56-70
15W- 5Mo- 1Zr- 0"1C 15W- 5Mo - 5Ti- 1Zr 10Mo - 1.5Zr - 0.3C - 0-03Ce,La (22 + 28)W- 2Hf-(0"067 + 0"13)C 30Hf- 9W a~, MPa
llO0°C 350-450
1200°C
5Mo - 5V- 1Zr 27Ta- 10W- 0.7Zr 4.6Mo - 1.4Zr- 0.12C 5W- 2Mo - 1Zr- (0"05 + 0.07)C 10W- 2"5Zr a~,MPa
1100°C 260-450
.
1200°C 180-350
Ol00,MPa 1300°C 140-300
llO0°C 210-300
1200°C 120-220
1300°C 100
*Annealed state. operating properties, to be used in new generations of spacecraft, aircraft and powder engineering, as well as in connection with development of new progressive technologies. In this search, a problem appeared in design materials intended for filling two operating temperature ranges. The first temperature range extends from 600 to 900°C where wrought Ti-alloys do not already work whereas Ni-alloys do not meet the specific strength requirements. The second one extends from 1100 to 1500°C where wrought Ni-alloys do not already work whereas alloys based on intermetallics and heat-resistant ceramics, considered as perspective structural materials, do not provide the required strength, ductility or fracture toughness. The present report, drawn up on the basis of further research carried out by the authors and the analysis of the published data, deals with the scientific principles of alloying and control of the structure of niobium alloys providing high strength and workability of the alloys for operating temperature ranges from 1100 to 1400°C and from 600 to 800°C. To design Nb-alloys destined to operate at temperatures where 0-1 Tin,it< T< (0"3 + 0"65) Tmelt all
the known methods of strengthening may be applied, such as the solid solution one, alloying by substitutional and interstitial elements, precipitation or dispersion strengthening, as well as creation of special structures. As a rule, the most effective combination of strength and ductility is achieved when joint application of several reasonably chosen methods of strengthening is used.
SOLID S O L U T I O N S T R E N G T H E N I N G Proceeding from the known mechanisms of solid solution strengthening, producing and some other requirements to be met by the designed alloys, the alloying elements should comply with one of the following criteria: (1) presence of wide solubility range even if at high temperatures; (2) an increase of intra-atomic bonds that is sufficiently reflected in increased melting point (Tme~t) of the alloy-strengthening in this case is kept at T> 0"5 Tmelt where the unstrengthening processes are determined by diffusivity of the elements; (3) a great difference of atomic sizes as compared to N b - this parameter determines the origin of considerable elastic distortions in the crystal lattice causing
Niobium--base alloys dislocation impeding the plastic flow process-such strengthening is kept at T< 0"5 Tmelt where the strengthening processes do not depend on diffusivity of alloying elements; (4) the alloy workability should not be radically decreased; (5) the alloy should not become considerably heavier (this is of great importance with respect to flying and rotating constructions) and should not be very expensive. W, Mo, Ta, and Re meet the requirements for operating temperature higher than 0"5 Tmelt.Content of these elements should not exceed 10-12 at.% to keep satisfactory workability of the alloys.3 A considerable difference of atomic sizes between Nb (Rat=l.46A) and Cr (1.256), V(1.31), Zr(1.59) and Hf(l'60) should provide strengthening of Nb with these alloying elements at operating temperatures below 0.5 Tme~t. Os, Ir, Ru and Re, meeting some of the requirements mentioned above, have high density (22.61, 22.65, 12.45 and 21.03 g cm -3, respectively) and are very expensive, therefore they should not be used for alloying Nb. Interstitial solid solution strengthening of niobium (by C, N, O, B) is not promising at high temperatures due to high diffusivity of these elements under these conditions. Moreover they decrease low-temperature ductility of alloys. These approaches based on analysis of binary systems are simplified to a certain extent. It is more reliable to estimate joint influence exerted by several alloying elements, i.e. on the basis of multicomponent systems.
STRENGTHENING BY DISPERSOIDS
The other method of increasing heat resistance of Nb is connected with precipitate and/or dispersion strengthening. Some oxide, carbide, nitride and boride dispersions as well as intermetallic ones or a combination of these phases can be used in this case. The strengthening phase should meet the following requirements: (1) higher mechanical strength as compared with that of the matrix; (2) particles must have sufficient size to be nonshearable obstacles to dislocation motion; (3) sufficient volume content to diminish Orowan loop; and (4) stability of particles size at operating temperatures. It is determined by high thermodynamic characteristics of the phases, its interaction with the matrix (components of the phase should exhibit low solubility and low diffusivity in the matrix) and low values of the surface energy with matrix.
305
The systems with two phase equilibria between Nb (or its solid solution) and such phases as intermetallics (Os, Ir, Ru, Pd, Pt, A1, Fe, Co, Ni and Cr), solid solutions based on alloying elements (Zr, Hf) or oxides, carbides, nitrides and borides both formed by Nb and the alloying elements can be revealed on the base of binary and ternary Nb phase diagrams. Intermetallic compounds are formed at relatively high concentrations of alloying elements (A1, Cr and Fe) radically lowering the temperature of Nb melting. So at temperatures exceeding 0.45-0-5 TrueStintermetallic phase particle coagulation is accelerated owing to an increase of diffusivity of the alloying element and self-diffusion of Nb. 4 Thus the precipitate strengthening by intermetallics can be effective for Nb below 1000°C.
STRENGTHENING BY OXIDES, CARBIDES AND NITRIDES
Strengthening of Nb at temperatures exceeding 1000°C should be provided by thermodynamically stable oxides, carbides, nitrides and borides. The most thermodynamically stable are oxides, carbides and nitrides of metals of Periodic System group IVA (Ti, Zr, Hf and Th) as well as the highest oxides of metals of group IIIA (A1, Y, La, Ce and Lu) and monoxides of metals of group IIA (Ba, Sr, Mg and Ca). When choosing phase, one should take into account compatibility of its elements with Nb minding an alloy producing process. So alloys with oxides of metals of group IIIA which are poorly soluble in Nb both in solid and liquid states, can be produced via powder technology. In this case the phases should keep thermal stability during compacting and sintering operations. Oxides, carbides, nitrides and borides of group IVA metals are dissociated and dissolved in the melt of high valency Nb ÷5 and during solidification are crystallized as separate phases or as solid solution of the elements in Nb. The alloys with these phases can be produced by melting. The researches, 5-15 accumulation and systematizationl 2,13,16-20 of phase diagrams Nb-Melv-X (X=O, N, C, B) showed that they are characterized by two types of isothermal section. The first (Fig. l(a) and (b)) shows that Nb can be in binary eutectic equilibria with MewX. It means that the phase diagram Nb-Mew-(O,N) has a pseudobinary eutectic section (Fig. l(d)). Presence of the ternary equilibria area Nb-Nb2X-Me~vX in
306
E. N. Shefiel, O. A. Bannykh
carbon-containing systems prevents the formation of a completely pseudobinary section (Fig. l(c)) but within a certain concentration range, along the Nb-MeivC direction, two-phase equilibria occur between Nb and (Nb,Melv)C. The second is a solid solution of NbC and MetvC. Such a Nb-MewC section may be considered as a partial pseudobinary one. Only those alloys which belong to the pseudobinary systems may exhibit the highest strength. Based on the laws of interaction between the phase diagrams and 'composition-high-temperature strength' diagrams 1,zl the highest strength belongs to those alloys of composition close to the solubility limit of MeivX in Nb (the I group of alloys) and to those which have hypoeutectic composition (the II group). The heterophase structure of the first of them is formed as a result of a solid solution decomposition whereas precipitation strengthening effect is determined mainly by dispersoid size. The heterophase structure of the II group of alloys is formed mainly during crystallization whereas strengthening is provided by a considerable growth of phase amount determined by the eutectic composition. 1 1 phase N ] 2phases Co)A ~3phases //" ~ )2 Nb~MeN
Nb
C
Meiv Nb
Melv
(d) ............... I
N Nb
Fig. 1.
i
btI~[~---II~ I I I ~ Groupsof alloys ,I MeW Nb
MeivX
Isothermic (a)-(c) and pseudo-binary (d) section of phase diagrams Nb-Melv-X (X = O,N,C).
The analysis of the own and published experimental data on concentration dependence of the strength of Nb-ZrO2(HfO2), Nb-ZrN(HfN) and Nb-ZrC(HfC) alloys and comparison of these data with the corresponding phase diagrams made it possible to designate the concentration ranges of alloys of the I and II groups (Table 2) in pseudobinary systems Nb-ZrOz(HfO2), Nb-ZrN (HfN) as well as in the system Nb-ZrC(HfC) the binary equilibria in which is realized not in the whole range of concentrations. Oxide strengthening
Oxide-containing alloys can be produced by charging the oxides during electric arc melting 16,17,22 as a result of 'contamination' during manufacturing and testing of material 23,24 by powder technology 25 or as a result of internal oxidization. 26-31 In ageing Nb-l% Zr(Hf)-O alloys the highest strengthening is achieved after quenching from 1700°C and following ageing at 950°C (Fig. 2). This effect is the result of precipitation of dispersed plate monocline phase ZrO2(HfO2) which is coherent to the matrix and oriented by its length a l o n g ( 1 0 0 ) N b . 16,17,22 In N b - T i - O 25 alloys the highest strengthening is achieved after ageing at 1000°C due to T i O 2 particles of 60 A size and with average distance between particles equal to 350A, coherent to matrix. 25 In the internal oxidized N b - l % Z r alloys while ageing at 1100°C, Z r O 2 particles of 40-60 A precipitate. 29-31 The thermal stability of MewO2 particle sizes in the ageing alloys depends greatly on what oxide forming element is present, in excess, in solid solution. Elevated content of Zr(Hf) in the solid solution above the stoichiometric M e i v O 2 facilitates the oxide particles coagulation. The accelerated growth of particles at 1100°C in the alloys with excess Z r 16,17,24 (Table 3) arises from the presence of free Zr(Hf) controlling the speed of growth of the phase particles and from
Table 2. Content of the dispersoids providing the precipitation (I group) and the dispersion (II group) strengthening in Nb-Meiv-X (X = O,N,C) alloys
System
Nb-ZrOz(HfO2) Nb-ZrN(HfN) Nb-ZrC(HfC)
Optimal content of dispersoids in alloys of I group
Content of dispersoids in alloys of H group
(rot. %)
(vol. %)
0.5-0"6 for at.%Zr/at.%O > 0"5 2 for at.%Zr/at.%O -<0-5 5 3"5
1-2-2 5-15 3-17
307
Niobium--base alloys /•..._.-e _•/
,, ~
•f ° /
Table 4. The size of ZrO2 particles in internal oxidized N b - l % Z r alloy annealed at 1600°C versus the free oxygen concentration in solid solution 29-31
Tq = 1700°C t = 2h --
\
o .~ y_.
.,.-~
-2 ~ .~
Annealing time
Diameter of particles (A)
~h)
Free oxygen concentration (at. '%)
-4 2800"
1
2400"
10 30
Z~
0"16
0"5
1.7
100 170 230
100 160 190
90 120 160
> ~Z 2100 1200 16_18x ~ 1700
1300! 600
_~ 1000 ~ ~'~, ~ _
800 10130 1200 Temperature of ageing (°C)
~8 0 06001~
~
o [] • • A ©
2% ZrO2,int. oxid. 1.5% ZrO2, aged 1% ZrO2, aged 0.3% ZrO2, recryst. 0.3% ZrO2, aged 0.12% ZrO2, aged
~ 6 " - 6 ,
%
Fig. 2.
Effect of quenching and following ageing temperatures on hardness and variation of resistivity for two-phase N b - 1 wt.%ZrO2 alloys with 0-3 vol.% Z r O 2.
"2
20x 400 ~•
Table 3. The size of Z r O 2 particles in N b - l % Z r - O alloyst with (0"3 + 0.6 ) vol.% Z r O 2 versus the ageing conditions
Ageing Plate length, Plate thickness conditions (A) (A)
Cc)
cA)
950
5
30
950
20
50
950 100
60
1100 1200
Distance between particles
20 5
150 570
Several atom layers Several atom layers Several atom layers 20 30-50
--
\ 18x
200
3 29 x
"A.~ I I 800 1000 Test temperature (°C)
I 1200
Fig. 3. UTS and elongation for two-phase N b - l w t . % Z r - O alloys with various contents of ZrO2 versus test temperature.
-70-140 200 400-860
t A f t e r quenching from 1700°C.
the increase of oxygen diffusivity in Nb at 1100°C. This results in reducing the effect of precipitation strengthening at this temperature. An efficient suppression of oxide particle coagulation up to 1600°C has been obtained in internal oxidized alloys with excess oxygen content (related to MetvO2)in solid solution (Table 4). The authors 29-3~ explain this fact by dramatic decrease of equilibrium solubility of ZrO2 in the solid solution in presence of excess oxygen. In order to provide sufficient ductility of alloys free oxygen content should not exceed 0.2-0.3 at.%. The structure of these internal oxidized alloys is characterized by presence of 1 vol.% of highly dispersed (40-60 A) coherent to the matrix particles of f.c.c, modification ZrO2(HfO2) precipitated from the solid solution. High elasticity
modules and high dispersivity of this phase preserved up to high temperatures, provides a considerably greater increase of recrystallization temperature in the internal oxidized alloys (up to 1 8 0 0 ° C ) 26'27 a s compared wth alloys having similar amounts of phase but with excess free Zr in solid solution (122 5°C) ~7and a greater effect of precipitate strengthening (Fig. 3). Oxide strengthening increases the creeprupture strength and creep resistance of alloys16,17,23,24.27,28,32,33 greater than tensile strength (Figs 3 and 4). The creep resistance is extraordinarily sensitive to extremely low oxide dispersoids content.23,24,2s,34 Oxide strengthening alloys exhibit satisfactory ductility (~p,6>_10%, Tbr_<-50°C)within the temperature range 20-1300°C (Fig. 3). This can be explained by a low volume content of oxide particles ( < 1 vol.%), their high dispersivity and small grain size. Exceeding 1.5 vol.% of oxide phase in alloys leads to formation of large (2-3 pm) hypereutectic oxide particles and to a
308
E. N. Sheftel, O. A. Bannykh
10° I
Table 5. The size of particles of alloys with ZrN versus ageing conditions
o IO00*C = 70MPa ol
--,
•
Time of ageing (h)
2
10-1
Size of particles (.4) Temperature of ageing (*C) 900
1-3 20 (") 10"2
30-60
0
1200
1400
1500
30-90 120-140 300-500 200-350 70-120 200-270 150t 80-200 250-330 120-300 600-800
50 1-3
tDistance between particles.
10-3 0
1000
I 0.3
I 1.0 Vol%Z r O 2
I 1.5
Fig. 4. Creep rate of two-phase Nb-lwt.%Zr-O alloys versus volume content of ZrO2: 1--recrystallize annealing; 2--quenching 1700"C + ageing 950"C, 10 h.
7(
great loss of strength. In this case the alloys are greatly embrittled. Nitride s t r e n g t h e n i n g
o
5(
r~
Nitride containing alloys can be produced by arc melting of nitride containing charge 16,17,35 or by melting in nitrogen containing plasma, 36 as well as by the diffusion saturation method 37 and powder process. Investigations 38-4° of ageing in Nb-ZrN(HfN) alloys have showed that active solid solution decomposition occurs within the temperature range 900-1200°C and leads to formation of 5-6vo1.% of highly dispersed f.c.c. ZrN(HfN) particles* (Table 2). The greatest effect of strengthening is achieved after quenching from 1700°C and following ageing at 1000°C. The precipitates of ZrN(HtN) after this condition in this case are of 150-450 A, the distribution density is (5-7)× 1016 particles cm -3, coherent with the matrix. Z r N ( H ~ ) particles stay highly dispersed up to 1500°C (Table 5). Thermal stability of a two-phase structure at 1000-1500°C exerts influence upon development of recrystallization processes occurring in the deformed alloys and taking place within the same temperature range. Highly dispersed nitrides (Table 5) slow down transformation of deformed :~Owingto formation of solid solutions between MewN and NbN in Nb-Meiv-N systems (Fig. l(a)) nitride (MewNb)N should be formed, however, presence of Nb in the nitride has not been determined experimentally probably because of its low content.
31
800
1000 1200 Test temperature(*C)
Fig. 5. UTS divided by alloy density for some commecial Nb-base alloys and for two-phase Nb-Zr-N alloys with various volume contents of phase ZrN versus test temperature.
structure and completion of the primary recrystallization up to 1400-1500°C. The nitride phase effectively slows down grain growth. The grain size in the alloys with 5 vol.% nitride is less than 15-20/~m after annealing at 1700-18000C for 1 h. 41 Tensile strength and rupture strength of the nitride containing alloys is increased within the temperature range 20-1200°C with an increase of the phase amount and its dispersivity 12,13,3s,39,42 and depends on heat treatment conditions.12,13,16,17, 37,42 Specified heat treatment gives these alloys strength and specific strength larger than that of many high-strength Nb-alloys with highly alloyed solid solution (Fig. 5, Table 6).
Niobium--base alloys It is known that the phase composition of the alloys obtained under real conditions may not comply with the stable equilibria phase diagram. This may be connected with stability of the metastable state or with deviation from the specified chemical composition. It is by these reasons that one can explain occurrence of large eutectic nitrides up to 3000/k in the structure of precipitation strengthened alloys. In these alloys, besides ZrN(HfN), hexagonal phase N b 2 N 3 8 - 4 ° having plate shape with average linear dimensions of (2000 + 10 000)/kx (300 + 1200) ,/k is released at the earlier stages of ageing at 800 to 1200°C. Nitride Nb2N is not stable and it dissolves in 2-5 h after its precipitation. Neither eutectic nitride or Nb2 N contributes to strengthening at 900-1200°C. Information about the alloys of the II group is limited by the alloy with 12vol.%ZrN. 14 The eutectic type structure of this alloy is represented by ZrN fibres 0-8-1 #m in thickness and up to 20 #m in length. Tensile strength and rupture
Table 6. R u p t u r e stress (o~00) of Nb-alloys with nitride strengthening
Alloy
Oloo, MPa Annealing at 1400°C
strength of this alloy is lower than that of the alloy with 4 vol.% ZrN up to 1300°C.
Carbide strengthening A great number of Nb-alloys developed in the 1960s and 1970s are based on precipitation strengthening by carbide (Table 1). The fact that at a Zr content within 1-2 mass% and C, 0.12-0.13 mass% (-1-1-5vo1.% of carbide phase) alloys characterized by satisfactory workability have predetermined compositions of the alloys developed in that time (Table 1). Formation of two-phase structure in the alloys with 1-2% carbide during cooling from the singlephase state proceeds with precipitation of the hexagonal carbide N b z C ( o r with Nb3C 2 at rapid cooling) in form of plates of 0.5-1 #m size. 15,43-45 It is not possible to suppress precipitation of these carbides even at such rates of cooling as ( 5 - 7 ) X 106 K s - l . 46 Nb2C is not stable at high temperature, while being heated up to 1000-1200°C it is transformed into stable f.c.c. carbide containing Mew, with a lattice parameter close to that of NbC (4-47 A). Completeness of transformation Nb2C to f.c.c, carbide as well as Zr or Hf content in it depends on ageing temperature a n d d u r a t i o n . 43,47
Quenching from 1700°C + ageing at IO00°C, 20 h
Test temperature (°C)
Nb- lvol.%ZrN Nb-2-5vol.%ZrN Nb-4vol.%ZrN Nb- 10%Mo-4vol.%ZrNt N b - 1-5vol.%HfN Nb-4"5vol.%HfN Nb-I 0%Mo-4vol.%HfNt
1100
1200
1100
1200
-120 ------
-65 ------
160 240 255 -----
90 110 130 140 100 145 185
? Q u e n c h i n g f r o m 1800°C + ageing at 1000°C, 20 h.
Phase transformations occurring in alloy with 1 mass% Zr and 0.12 mass% C were investigated firstly in Refs 43 and 47. Developed heat treatment conditions have provided increasing the strength of such alloys in ever higher degrees (Table 7). The alloys with 1-1.5vo1.% carbide keep their strength up to 1100-1200 ° (Fig. 6). Information concerning alloys with carbide phase content higher than 1-2 vol.% is limited by the data concerning alloys VN-4; 1 P W C - 3 3 49 a s well as by the results of analyses made by the authors j0-52 Arc melted alloys containing up to 12 vol.% carbide dispersoids have been investigated in Refs
Table 7. M e c h a n i c a l p r o p e r t i e s o f N b - 1 0 % W - 5 % M o - l % Z r - 0 - 1 2 % C
Treatment
Forging Quenching f r o m 1700°C Quenching + ageing LTMTt
309
alloy after various t r e a t m e n t s (by O. B. Dashevskaya)
Mechanical properties at 1100°C oB (MPa)
ao.2 (MPa)
6 (%)
-46
43 28
-31
44
29
35
68
56
17
t L o w - t e m p e r a t u r e thermo-mechanical treatment.
Oloo (MPa) 24 --
28.5
Life-time (h) 60 1200
310
E. N. She#el, O. A. Bannykh 500 Nb - 15W - 5Mo - 1Zr - 0,1C (F-48) 450 -
%-
~4ooe~
350 -
300 o
E .=
250 -
200-
150 1050
I
I
I
I
I
1100
1150
1200
1250
1300
Test temperature (°C)
140
¢•-
Fig. 7. Microstructure of the Nb-7vol.%(Nb,Zr)C alloy. Optic micrographs: (a) as cast Vcoot=10 ~ K s-l; (b) after hydrostatic extrusion with counter-pressure, 25% reduction; (c) see (b), 60% reduction; (d) ribbon of 150 # m thick, Vcool= 105 K s-l; (e) transmission electron micrograph, see
10W - 1Zr - 0.1C (D-43)
120
1~--
(d).
80--
Table 8. Size of carbides (Nb,Zr)C particles in Nb alloy with 6 vol.% carbide versus annealing conditions
Nb - 10W - 0.15C
20 1050
I
I
I
I
I
1100
1150
1200
1250
1300
Test temperature (°C)
Fig. 6.
Annealing conditions
(°c)
Mode of particles
(h) Eutectic,
(#m)
Result of dispersion strengthening of Nb alloys by 1.2 vol.% carbide phase (Nb,Zr)C.
50-52. The phase composition of the alloys is represented by Nb-base solid solution and f.c.c. carbide (Nb,Zr)C being in eutectic equilibria. Their as-cast structure is represented by large, 5 - 1 0 # m , nonstoichiometric eutectic carbides (Nb,Zr)xC1-x with a lattice parameter close to that of NbC (4.47 A). Zr concentration in the carbide and a degree of its approximation to stoichiometric composition increase with increasing Zr and C content in the alloy. Being precipitated primarily at grain boundaries, these carbides form an almost continuous network along the boundaries (Fig. 7(a)). Annealing of cast alloys leads solid solution decomposition and precipitation of (Nb,Zr) C particles oriented by their length along {100} Nb (Table 8). Alloys with 5-12vo1.% carbide have much higher strength and specific strength up to 1400°C in comparison with those of alloys containing 1-2% of the phase (Fig. 8). They have high
Particle size
1000 1200 1200 1200 1200 1600 1800 2100 2100
1 1 5 20 50 1 1 1 1
2-8t 2-8t 2-8 2-8 2-8 2-8 2-8 up to 6 up to 6
Segregatedfrom solid solution Plate length (,4)
Platewidth (.4)
100-150 600-800 1000-1500 1300-1700 2300-3400 1900-2500 5000 ---
20-30 50-100 200-300 200-400 350-400 300-400 500-600 ---
tit depends on crystallization rate.
capacity for strain hardening and high recrystallization temperature (1500°C). However, a disadvantage of these alloys is their poor ductility and fracture toughness at room and medium temperatures, which prevents their application. One of the main reasons for it is large carbide particles making up a continuous network along grain boundaries (Fig. 7(a)). During plastic deformation local stress concentration at large carbide particles promotes local fracture away from the matrix (which occurs more easily at matrix-particle
Niobium--base alloys × ",~ • I zx 120"" • " [ - xD
Mo, 5vo1% HfN, 80/~ 1.5% ZrO2, int. ox., 100/~ Mo 4% ZrN, 80/~ 4%'ZrN, 80/~ 6% (Nb Zr) C h-extr. (2000 x 300) ,~ Mo, 6%'(Nb, Zr) C, cast (4000 x 400) .~
1000 --o x~
g ~o~ 52" .-~.."-~,L--a--..×sx
:..;&5,,,\%on
600
- - ..--
-...
o~ o
NX×
11 x
- .,2,X~xl2X ~ ~i5 x
400
if-
~ 2000
- " " ~
r~37x
~ ~ " ~ " - - -'~e~N• 1.4% Zr02, 5500A " ~ ' ~ ' ~ - ~ . o 0.3% ZrO2, 40A • O 1,~% (Nb, ~r)C, (1300 x 250)i~ 700
Fig. 8.
~.
Comparison
900 1100 1300 Test temperature (°C)
1500
o f oxide, n i t r i d e a n d c a r b i d e s t r e n g t h e n i n g effect.
interface). As a result, the cracks nucleate at these sites. 5° Increase of low-temperature ductility of cast alloys with high carbide content can be achieved via optimization of cast structure by means of proper heat treatment, high-temperature extrusion (1700-1800°C) or by applying such processes as cold hydroextrusion with counterpressure ~7.50and rapid solidification.5°-52 Multitransitional hydroextrusion with counterpressure (30-15 kbar) made it possible to deform the alloys with 5-7 vol.% of the phase with reduction of up to 80-90%. The deformation leads to breaking of large eutectic carbides, distintegration of the grain boundary continuous carbide network and to increasing of structure uniformity (Fig. 7(b) and (c)). The hydroextrusion leads to the formation of a preferred texture of (110) type in the extrusion direction. The above factors improve material ductility. Along with the increase of ductility, hydroextrusion results in a considerable strain hardening that is retained up to 1500°C. Rapid solidification of the alloys with 5-9vo1.% carbide, carried out by different methods providing a cooling rate of 105 K s -1, leads to refining of the structure (Fig. 7(d) and (e)). The grain size after rapid solidifcation is 1-3/~m, and that of eutectic carbides--0.2-0.3 ~m at decreasing of their total amount and disappearance of grain boundary continuous carbide network (Fig. 7(d) and (e)).
311
All the above described should provide an increase of strength (according to Hall-Petch) and ductility. Increasing ductility is provided by facilitating propagation of slip from one grain to an adjacent one across the 'feed' boundary and decreasing the contribution from matrix-particle interfaces along which the nucleation of cracks occur. Annealing of rapidly solidified alloys at temperatures up to 1800°C does not lead to grain growth and coagulation of eutectic carbides but leads to decomposition of the solid solution with formation of the dispersed (Nb,Zr) C (Table 8). The above results allow us to suppose that the material produced via rapid solidification followed by compacting would exhibit a unique combination of strength and ductility. The alloys of eutectic composition produced by directed crystallization deserve attention. Directionally crystallized Nb-Nb2C alloy containing up to 30vo1.% carbide as fibres have the specific strength of 4.5 km at 1200°C. 53 Directionally crystallized Nb-(Nb,Mew)C eutectic alloy possesses the specific strength of 6 km at 1200°C exceeding that of all known Nb-alloys. Comparison of oxide, nitride and carbide strengthening Comparison of the above, within 20-1400°C, has been made taking into account superposition of such factors as thermodynamic stability, quantity and dispersivity of the strengthening phases (Fig. 8). Thermodynamic stability of the phases increases in the order: carbides, nitrides, oxides. The free energies (AH~98,kJ mol-1) a r e : N b C 140"3; Z r C - 200; H f C - 230.1; Z r N - 365"3; HfN - 369"2; ZrO 2 - 1094.1; HfO 2 - 1113. The amount of strengthening phase is increased from oxides to carbides and then to nitrides for I group alloys and from oxides to nitrides and carbides for II group alloys (Table 2). In I group alloys the highest strengthening up to 1200-1300°C is provided by nitrides formed in the alloys after quenching followed by ageing in the greatest amount (4-5 vol.%) and the highly dispersed state (70-90 A). The alloys with 1-2.5vo1.% carbides exhibit the lowest effect of strengthening. The cause of this is the low thermodynamic stability of the phase formed in these alloys during ageing (NbC) and the relatively large size of its particles. Strength of precipitate strengthened alloys with -< 1.5 vol.% ZrO2(HfO2) is intermediate between those of carbide and nitride containing alloys. But
312
E. N. Sheftel, O. A. Bannykh
the strength can be increased up to that of nitride containing alloy by providing small size ( < 100 A) oxide particles up to 1600°C due to an excess content of oxygen (related to ZrO2(HfO2)) in solid solution (Fig. 8). The strength of alloys with 1-2vo1.% (Nb,Melv)C is inferior to that of alloys with oxides and nitrides, but slower with temperature increasing degradates. The latter is provided by lower solubility of carbides in Nb matrix and slower coagulation of multicomponent phase (Nb,Mew)C particles by comparison to two-component oxide or nitride particles. The alloys with 5-9 vol.% carbide dispersoids exhibit greater high-temperature strength than alloys with 1-2 vol.% carbide, but are inferior to precipitate strengthened alloys with nitride. The strengthening by phases with Hf is more effective than that by phases with Zr. JOINT SOLID SOLUTION AND OXIDE, NITRIDE AND CARBIDE STRENGTHENING Alloying of Nb by W, Mo, Ta and V, as well as by Melv and one of the interstitial elements (O,N,C) provides joint solid solution and dispersoid strengthening at 20 to 1400°C (Fig. 9). The contribution from each of them to the general strength level depends on alloy chemical composition, phase content and testing (operating) temperature. Thus, at 1300-1400°C (0-55-0.6 Tmelt) when solid solution strengthening in Nb-alloys is being decreased, the contribution of dispersion strengthening by carbide and nitride becomes important.
900[ - - ~ 800~ ~
~= 700l--
UTS attainable by MelvC,MeivN ~
/
~ 600q) 500'x'- ~
Two-phas,~ N
* 400 --~1 Single-phase~ | "~ ~~ ,.-- ~ alloys 300-I"
__,~, 600
Fig. 9.
I
~/~///A~ W(,ff//~7/fl~
I
1000 1200 Test temperature(*C)
1400
Ultimate tensile strength of single-phase and twophase Nb-alloys.
Solid solution alloying by W of Nb-MewX alloys leads to a decrease of their low-temperature plasticity, 12,13 while alloying by Mo leads to an increase. 54 Simultaneous increase of tensile strength, stress-rupture strength and plasticity of alloys with Mo (Fig. 8) are provided by solid solution strengthening and Mo influence on the twophase structure formed in the alloys. The influence has been carried out by following the refining of grain size, decrease of interstitial element content in solid solution providing increase of precipitating phase amount as well as increase of strength of particle-matrix interface. PRECIPITATON STRENGTHENING BY INTERMETALLIC PHASES Increasing demands for new materials used in aircraft engine construction, conditioned by the elevation of operating temperatures of the engine, stimulate a search for the materials which can fill the operating temperature gap from 600 to 800°C. At these temperatures, high temperature strength Ti-alloys can not work whereas Ni-alloys do not meet the specific strength requirements. Russian scientists have developed high-temperature, heat-resistant, wrought alloy, operable at 600-700°C. The principle of its alloying lies in the choice of elements providing formation of protective films which increase resistance of the alloys against oxidization (Ti,AI) and formation of precipitate strengthening phase (Ti3A1). The specific yield strength of the alloy is considerably higher than that of Ni-high-temperature wrought alloy operable at 600-800°C. When stress-rupture tested in air, the alloy exhibits significant stress riser (notch) sensitivity at 600-700°C. Moreover, specific stress-rupture strength of the alloy becomes considerably lower than that of Nialloys. The cause of this phenomenon is the oxidizing effect of the environment. 55 The notch sensitivity was shown to be controlled by oxygen that is absorbed from the environment to form a special underscale layer of altered chemistry and phase constitution. The environmental damage could be prevented to a large extent by optimization of the structure and alloying with carbon. A perspective of heat-resistance increasing of high-temperature alloys of the Nb-Ti-Al system may be their additional alloying by Cr. The dispersed intermetallic phase NbCr 2 is expected to be more stable against coagulation in Nb-matrix than Ti3AI and therefore to provide more effec-
Niobium--base alloys 1000 _-800 v
(VM4ArrAN)
600 --
~'5oo
_
~
\
REFERENCES
Mo-alloys (TzM-5, TzM-6)
~
~
v ~ L N - 1 TMT . ~_ 5VMTz
] 1
LN - ~
!
~
~ i n t . oxid. f ~-(~o, rio C
directed eutectic
/ e 200- (IN-100)-/~x ~
"~////_/(/
Ni-~loys, directed ~ V ~ eutectic ~ ~ .
~
\\'v~(g,f~//~
_
,
._ \ _ X
-
-
4
".
Nb-9Mo-4ZrN D-43
o ---~, 700
I
900
I
I
1000 1100 Test temperature (°C)
I
1200
I
1300
Fig. 10. The 100 h stress rupture strength of Nb-base alloys in comparison with high-strength Ni- and Mo-base alloys.
tive strengthening. Cr is able to increase the heatresistance of Nb-Ti-AI alloys by 100°C to 800°C, due not only to Cr-containing protective surface film formation, but also to holding of Nb in NbCr2, decreasing the amount of Nb moving towards the surface layer.56
CONCLUSION The above described data and concepts can be the basis for the most efficient choice of Nb-base alloy compositions, methods and conditions of production and treatment, providing high-temperature strength, workability and, for medium operating temperatures, high heat-resistance. This approach allows development of Nballoys with high-temperature strength, higher than that of Ni-superalloys above 800°C and similar or even higher than that of Mo-based alloys at 1000-1300°C (Fig. 10). This group of Nb-alloys requires special coating protecting against oxidization. The approach also provides for development of Nb-alloys to be used at medium temperature, being heat resistant, with a specific strength at 600-800°C higher than that of high-temperature wrought Ni-alloys.
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